Cubic (zinc-blende) aluminum nitride

ABSTRACT

Device quality, single crystal film of cubic zinc-blend aluminum nitride (AlN) is deposited on a cubic substrate, such as a silicon (100) wafer by plasma source molecular beam epitaxy (PSMBE). The metastable zinc-blend form of AlN is deposited on the substrate at a low temperature by a low energy plasma beam of high-energy activated aluminum ions and nitrogen ion species produced in a molecular beam epitaxy system by applying a pulsed d.c. power to a hollow cathode source. In this manner, films having a thickness of at least 800 Å were produced. The lattice parameter of as-deposited films was calculated to be approximately 4.373 Å which corresponds closely to the theoretical calculation (4.38 Å) for cubic zinc-blend AlN. An interfacial layer of silicon carbide, specifically the cubic 3C—SiC polytype, interposed between the epitaxial film of zinc-blend AlN and the Si(100) wafer provides a template for growth and a good lattice match. The epitaxial layer of zinc-blend AlN has been characterized for its physical and optical properties. As a result, experimental data confirmed that zinc-blend AlN is an indirect semiconductor and has a bandgap about 5.34 eV. Due to the extraordinary piezoelectric properties of zinc-blend AlN, an illustrative device embodiment is a surface acoustic wave (SAW) device comprising interdigitated electrodes deposited by conventional means on the surface of the epitaxial layer of zinc-blend AlN to convert an electrical signal to a surface acoustic wave and vice versa.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application claims the benefit of provisional patent applicationNo. 60/128,662 filed Apr. 18, 1999 and assigned to the assignee herein.

BACKGROUND OF THE INVENTION

1. Field of the Invention

This invention relates generally to aluminum nitride (AlN) and moreparticularly, to epitaxial cubic (zinc-blende) AlN films that may have athickness on the order of 1000 Å or greater and a method of making sameby plasma source molecular beam epitaxy (PSMBE).

2. Description of the Related Art

The Group III-V nitride semiconductors (GaN, AlN, and InN) are of greatinterest for their potential as optoelectronic materials. Thesematerials have an equilibrium crystal structure which is wurtzite, orhexagonal. The bandgaps of the wurtzite nitride semiconductors are alldirect and their alloys have a continuous range of direct bandgapsvalues ranging from 1.9 eV for InN to 4.0 eV for GaN to 6.2 eV for AlN.As optical materials, these semiconductors are active from the orangeinto the ultraviolet.

Formation of nitride semiconductors for device applications requires,among other things, achieving the correct stoichiometry, inducing thecorrect energy to form a highly crystalline matrix, maintaining highpurity, and matching the lattice parameters of the semiconductor and thesubstrate. Much effort was expended in the 1960's and 1970's to grow andcharacterize Group III-V nitride semiconductors. However, the effort wasineffectual to achieve high-quality material. Recently, there has beenrenewed effort to create higher quality Group III-V nitridesemiconductors. However, GaN, AlN, and InN produced by conventionalmethods have high n-type background carrier concentrations resultingfrom native defects commonly thought to be nitrogen vacancies. Nitrogenvacancies affect the electrical and optical properties of the film.Oxygen contamination is also a major problem. Thin layers of AlN havebeen prepared by magnetron sputtering, chemical vapor deposition, ionbeam sputtering, and ion beam assisted deposition. However, thesemethods operate at elevated temperatures and generally do not result inepitaxial growth (i.e., growth oriented in one direction). Moreover,while these techniques have been successful in producing polycrystallineAlN films, they have not been successful in producing electronic-gradesingle crystal films.

AlN, in particular, is a promising material for high-power,high-temperature optoelectronic devices since it has very high chemicaland thermal stability, good thermal conductivity, and fast Rayleighvelocity. AlN crystallizes, under normal conditions, into thethermodynamically stable hexagonal wurtzite structure. However, themetastable cubic zinc-blende structure is expected to be easier to dopeand to have decreased phonon scattering, and therefore, to have higherballistic electron velocities, thermal conductivity, and acousticvelocities due to its higher symmetry. These properties give rise tomany exciting potential device applications.

There have been several reports of AlN having the metastable cubiczinc-blende structure. These reports, however, lack detail on thephysical, electrical, and optical properties of cubic zinc-blende AlNbecause the films were too thin for such studies, and certainly too thinto be useful for optoelectronic devices which require thicknesses on theorder of at least 2000 Å, and preferably 4000 Å to 8000 Å. The latticeconstant of zinc-blende AlN was calculated theoretically to be 4.38 Åusing data from the elastic constants of wurtzite AlN. This value waslater confirmed experimentally on a 12 nm thick film of zinc-blende AlNgrown pseudomorphically on cubic TiN sandwiched between a tetragonalAl₃Ti overlayer. To date, however, there have been no reports ofsuccessful fabrication of thick, device-quality films of zinc-blendeAlN. The known AlN films have been mixed hexagonal and possibly cubic(which could be the rock salt structure).

It is, therefore, an object of the invention to prepare zinc-blende AlNof sufficient quality and thickness to characterize it for itsmechanical, optical, and electrical properties and to be useful fordevice fabrication.

It is also an object of the invention to prepare device quality, singlecrystal, epitaxial films of cubic zinc-blende AlN.

It is a further object of the invention to produce a semiconductordevices that include an epitaxial film(s) of single crystal zinc-blendeAlN.

It is an additional object of this invention to provide a method ofmaking an epitaxial film of zinc-blende AlN.

SUMMARY OF THE INVENTION

The foregoing and other objects, features and advantages are achieved bythis invention which is, in a first device embodiment, a film of devicequality, single crystal cubic zinc-blende AlN. In other embodiments, thezinc-blende AlN film is deposited on a substrate, and preferably on asubstrate having cubic symmetry on its surface, such as a silicon (100)wafer (Si(100)). In a particularly preferred embodiment, there is abuffer, or an interfacial, layer of cubic 3C—SiC between the epitaxialfilm of zinc-blende AlN and the substrate which may be a Si(100) wafer.

In a specific illustrative embodiment, a semiconductor device comprisesa Si(100)-oriented substrate; an interfacial layer of 3C—Si(C) on theSi(100) substrate; and a film of single crystal zinc-blende AlN having athickness of at least 800 Å, and preferably in the range of 1000 Å to2000 Å, which is epitaxial with respect to the Si(100) substrate. Theepitaxial relationship between film and substrate is (100)AlN∥(100)Siand [101]AlN∥[101]Si. The interfacial layer may have a thickness rangingfrom several atomic layers (e.g., about 25-30 nm) and up.

In accordance with the principles of the invention, the metastablezinc-blende form of AlN is deposited on the substrate by a plasma beamof aluminum ions and activated nitrogen ion species produced in amolecular beam epitaxy system by applying a pulsed d.c. power to ahollow cathode source. In this manner, films having a thickness of atleast 800 Å were produced. Thickness, of course, is a function ofdeposition time, and films ranging from 10 Å to several microns, arepossible by the method of the present invention. The lattice parameterof the as-deposited films was calculated to be approximately 4.373 Åwhich corresponds closely to the theoretical calculation (4.38 Å) forcubic zinc-blende AlN.

The zinc-blende AlN epilayer films of the present invention have a widebandgap (experimentally determined to be about 5.34 eV); thermalstability (up to about 800° C.), and extraordinary piezoelectricproperties. In addition to the foregoing, the films have been ofsufficient quality to enable experimental confirmation that zinc-blendeAlN is an indirect semiconductor. When characterized in situ byReflection High Energy Electron Diffraction (RHEED), the films showfour-fold symmetry rather than the six-fold symmetry which is typicalfor hexagonal AlN in (0001) orientation. Furthermore, the RHEED patternsappear to be very similar to those for the Si(001) substrates, exceptfor different streak spacings. X-ray diffraction (XRD) revealed broadpeaks at diffraction angle (2θ) values of approximately 41° and 89.8°.These peaks match the (002) and (004) peaks of zinc-blende AlN with alattice parameter of 4.38 Å. Transmission electron microscopy (TEM)confirmed that the AlN produced by the method of the present inventionis cubic, single crystal and epitaxial with respect to the Si(100)substrate.

In accordance with the invention, the growth surface of the substrateshould preferably have a cubic structure to act as a template for cubic(zinc-blende) growth. Specific illustrative examples include, but arenot limited to Si(100) and magnesium oxide (MgO (100)). Of course, thesubstrate may comprise one or more layers. Preferably, the growthsurface of the substrate should have an good lattice match with AlN. Agood lattice match is defined as being within about 1%. However, as isknown in the art, substrates with seemingly poor lattice matches (e.g,Si(100) has a 19% mismatch), may be used since epitaxially depositedlayers have mechanisms to compensate for the mismatch, such as byforming defects.

3C—SiC is an example of a substrate that has less than 1% latticemismatch with AlN. In a particularly preferred embodiment, there is abuffer layer, or an interfacial layer, of cubic 3C—SiC between theepitaxial film of zinc-blende AlN and a Si(100) wafer. The 3C—SiC layermay have a thickness ranging from several atomic layers (e.g., about25-30 nm) and up. A silicon substrate with a 3C—SiC layer can bepurchased from a commercial source or made in the laboratory. The 3C—SiClayer can be deposited on a silicon substrate by any number of known,high temperature processes, such as chemical vapor deposition orenhanced MBE. As will be described more completely hereinbelow, 3C—SiCcan be made in situ on a silicon substrate by the methods in accordancewith the invention herein.

In a specific illustrative embodiment of one of the many deviceapplications contemplated by the invention, an interdigitated transducersurface acoustic wave (SAW) device comprises a Si(100) substrate, anepitaxial film of single crystal, zinc-blende film of AlN deposited onthe substrate, and two interdigitated electrodes deposited on theepitaxial AlN film by standard photolithographic techniques.

In a method embodiment of the present invention, a single crystal,epitaxial AlN film having a zinc-blende cubic structure is formed byexposing a heated substrate to a low energy flux of target atoms in anultrahigh vacuum PSMBE system. The PSMBE system uses a plasma depositionsource which is a magnetically-enhanced, generally cylindrical hollowchamber comprising a cathode. The chamber is lined with the targetmaterial which, in the present case, is MBE-grade aluminum. The targetmaterial is milled so that its thickness is greater at the upper, orexit, end of the chamber than the thickness at the lower end. In apreferred embodiment of the invention, there is about a 3° internaltaper in the chamber. Plasma is formed in the chamber by the applicationof d.c. or r.f. power. The application of a pulsed d.c. power producesan epitaxial layer of metastable zinc-blende AlN. The magnetic field andthe taper of the interior of the cathode cooperate to confine the plasmato the cathode. The low energy flux of target atoms is extracted fromthe exit end of the chamber either by the action of an impellerrotatably mounted in the cathode source or by an acceleration biasapplied to the substrate.

Secondary electrons are confined to the hollow cathode source and do notinteract with the substrate which is mounted distant from the exit endof the chamber by, in a specific preferred embodiment, approximately 25cm. Due to the extreme anisotrophy of the kinetically ejected atomsperpendicular to the source wall and the tapered geometry of the hollowcathode chamber, virtually no high-energy atoms or ions are directed tothe substrate surface. Instead, the atoms go through multiple collisionsand thermalize. Since ions are extracted from the source via an impelleror acceleration bias, the energy distribution is controlled. Massspectrometry energy analysis of the non-accelerated ions ejected fromthe r.f. powered source indicates a gaussian energy distribution aboutthe approximate 1 eV range.

The PSMBE system of the present invention, in contrast to planarsputtering systems, allows enough adatom energy to create AlN crystals(or other semiconductor nitrides) while minimizing damage to theunderlying crystal substrate. When pulsed d.c. power is used instead ofr.f.power, the prolonged high potential state of the pulse combined withthe delay time between pulses results in more energetic plasma with highplasma and Al/N kinetic energies. The increased intensity of the plasmagives more dissociated and hence, more energetic, aluminum and nitrogenspecies for deposition on the substrate. This increase in the plasmaenergy and active energetic species assists in the formation of cubiczinc-blende AlN.

In a further method aspect of the present invention, the substrates arepre-treated to de-grease the substrate with solvents and to removesurface oxidation. The substrate is then, preferably pre-heated to ahigh temperature, illustratively approximately 800° C., for one hourbefore deposition of the epixtaxial film.

In a specific illustrative embodiment, a Si(100) substrate ispre-treated by (i) ultrasonic cleaning in acetone for 20 minutes; (ii)blow-drying with dry nitrogen; (iii) ultrasonic cleaning in methanol for20 minutes; and (iv) blow-drying with dry nitrogen. The substrate isthen etched in an acid, illustratively 10% (by volume) HF to remove theSiO₂ layer from the surface and to saturate dangling Si bonds withhydrogen atoms. The substrate is then loaded into the depositionchamber, preferably within 10 minutes of treatment.

In a particularly preferred embodiment, the pre-treated Si(100)substrate is subjected to a source of carbon to combine with the Si,during the pre-heat, to create an interfacial layer of cubic 3C—SiC onthe Si(100) of at least several atomic layers thickness. The cubic3C—SiC acts as a template for the deposition of cubic zinc-blended AlN.Illustratively a source of carbon can be introduced during the etchingstep by including trace amounts of a mixture of hydrocarbons in theetching solution. For example, the hydrocarbons may be the resins foundin standard photoresist materials.

In an additional embodiment of the invention, a PSMBE system useful inpreparing epitaxial films on a substrate comprises an ultrahigh vacuumgrowth chamber. An ultrahigh vacuum pump is connected to the growthchamber so that base pressures in the range 10⁻⁹-10⁻¹⁰ torr can beachieved and maintained.

A substrate holder is, preferably, rotatably mounted in the interiorupper end of the growth chamber. The substrate holder is heated to bringa substrate mounted thereon to growth, or processing, temperature. Asubstrate bias power supply is electrically connected to the substrateholder.

At least one plasma deposition source, and preferably from two to eightsources, is located distal to the substrate holder to achieve anappropriate distance, preferably in the range of 5 cm to 50 cm, from theexit end of the source to the deposition, or growth, surface of thesubstrate. The plasma deposition source, in a preferred embodiment, is amagnetically-enhanced hollow cylindrical cathode lined with targetmaterial which, in this specific embodiment, is aluminum. The plasmadeposition source includes an inlet for high purity process gases,specifically nitrogen and argon. A source of pulsed d.c. or r.f. poweris electrically connected to the plasma deposition source for creating aof plasma of high-energy aluminum, nitrogen, and argon species. Animpeller is rotatably mounted in the plasma deposition source forejecting ions from the source. The cathode source is surrounded by ananode, which may be a stainless steel enclosure, at an insulatingdistance from the cathode.

In a preferred embodiment of the invention, the plasma deposition sourceis a cathode comprising a hollow, generally cylindrical chamber, thechamber having an opening at the upper end for emitting the generatedplasma and an inlet at the lower end for the process gases. The sourceincludes magnetic field generating means. In this specific embodiment,the magnetic field generating means is an array of magnets embedded inaluminum and forming the exterior wall of the chamber. The magneticfield confines the plasma to the hollow chamber. The target materialforms the interior lining of the chamber wall. The thickness of thetarget material lining the chamber increases from the lower end to theupper end to provide a taper that further confines the plasma to theinterior of the chamber. In a specific embodiment, the lining has a 3°taper.

BRIEF DESCRIPTION OF THE DRAWING

Comprehension of the invention is facilitated by reading the followingdetailed description, in conjunction with the annexed drawing, in which:

FIG. 1 is an schematic cross-sectional representation of a bilayersemiconductor device in accordance with the present invention;

FIG. 2 is an illustrative schematic diagram of a PSMBE system useful inthe practice of the invention;

FIG. 3 is a schematic representation of a magnetically-enhanced hollowcathode plasma deposition source in accordance with the presentinvention;

FIG. 4 is a RHEED pattern from a zinc-blende AlN film fabricated inaccordance with the method of the present invention, taken at azimuthsof [101], [100], and [310];

FIG. 5 is a deflection image taken by atomic force microscopy (AFM) ofthe surface of an AlN film fabricated in accordance with the presentinvention, a height image is shown in the inset;

FIG. 6 is a TEM image of a pyramidal pit of the type shown in the AFMimages of FIG. 5;

FIG. 7 is an Auger Electron Spectroscopy (AES) depth profile onzinc-blende AlN film specimens (FIG. 7a and FIG. 7b) fabricated inaccordance with the present invention and a non-zinc-blende AlN film(FIG. 7c);

FIG. 8 is a High Resolution TEM (HRTEM) lattice fringe image in the[110] orientation of a specimen having a zinc-blende AlN film on aSi(100) substrate showing an interfacial layer at the zinc-blendeAlN/Si(100) interface;

FIG. 9 shows symmetric XRD scans of two specimens of AlN grown onSi(001) substrates, an inset shows rocking scans of the same twospecimens;

FIG. 10 is a cross-sectional view of the AlN/Si(001) interface in the[1{overscore (1)}0] orientation taken by transmission electronmicroscopy;

FIG. 11 is plan view TEM selected area diffraction (SAD) pattern in the[001] orientation which demonstrates that the AlN is a cubic zinc-blendeAlN (β-AlN) structure;

FIG. 12 is a cross-sectional SAD pattern of the AlN/Si (100) interface;

FIGS. 13a and b are extinction coefficient (k′) data and refractiveindex (n′) data, obtained by spectrophotometric ellipsometry, from asample film of zinc-blende AlN on Si(100);

FIG. 14 is a plot of extinction coefficient data as (Eα′)^(½) as afunction of photon energy in eV that provides an experimental estimateof the bandgap of zinc-blende AlN; and

FIG. 15 is a schematic representation of an interdigitated transducerSAW device fabricated from the zinc-blende AlN of the present invention.

DETAILED DESCRIPTION

In an illustrative device embodiment 8, shown schematically incross-section in FIG. 1, a single crystal zinc-blende AlN film 5 isepitaxially deposited on a face surface of a substrate 6,illustratively, Si(100). In this specific illustrative embodiment, theAlN film is epitaxial with respect to the Si(100) substrate and has athickness on the order of 1000 Å. The epitaxial relationship betweenfilm and substrate is (100)AlN∥(100)Si and [101]AlN∥[101]Si.

In particularly preferred embodiments, a buffer layer of a materialhaving a cubic structure and a good lattice match with the zinc-blendeAlN film is interposed between epitaxial layer 5 and substrate 6. In aspecific illustrative embodiment, the buffer layer is an interfaciallayer of 3C—SiC. An interfacial layer of 3C—SiC is clearly shown in FIG.6 as reference numeral 7.

Devices in accordance with the present invention are preferablyfabricated by the PSMBE technique described in detail hereinbelow.

An illustrative schematic diagram of a PSMBE system useful in thepractice of the invention is shown in FIG. 2. An ultrahigh vacuum (UHV)growth chamber 10 is pumped by an ion pump 11 (e.g., 500 l/s) to achievebase pressures in the range 10⁻⁹-10⁻¹⁰ torr. In the specific embodimentshown, a throttled cryopump 12 assists in maintaining the vacuum underdynamic gas flow conditions. Although not specifically shown in thefigure, a sample load lock system maintains vacuum chamber integrity.The temperature of growth chamber 10 is monitored by a temperaturemeasurement system 13, such as an IR temperature measurement system.

The growth chamber may also be fitted with observation and analyticalequipment. In the embodiment shown, a differentially-pumped RHEED gun 14and RHEED screen with charge-coupled device (CCD) camera 15 areincluded. In addition, the embodiment of FIG. 2 has as an ellipsometrysystem 16 for evaluating optical properties.

A substrate holder 17 having a mounting surface 29 is rotatably mountedto the upper interior surface 18 of growth chamber 10. Substrates arepreferably rotated continuously during deposition to insure uniformthickness. Large substrates (up to 3 inch wafers) may be usedsuccessfully in this manner. Substrate holder 17 is heated,illustratively radiantly or resistively, to bring a substrate (notspecifically shown in the figure) mounted on surface 29 to hightemperatures, preferably between about 400° C. to 800° C. Substrateholder 17 is electrically connected to a substrate bias power supply 19.

There is at least one plasma deposition source, illustratively, plasmadeposition source 20 mounted to the lower interior surface 28 of growthchamber 10 at an appropriate target-to-substrate distance from the face,or growth, surface of a substrate mounted on substrate holder 17. Inthis specific embodiment, the distance is 25 cm. The source may bemounted off-axis up to 20° from normal. In accordance with the presentinvention, plasma deposition source 20 is a magnetically-enhanced hollowcathode plasma deposition source, lined with metallic target material,that is described in greater detail with respect to FIG. 3.

Process gases, specifically argon and nitrogen in this embodiment, fromsources 21 and 22 are purified to sub-ppb levels in titanium and coppergettering furnaces 23 and 24, respectively. The purified gases areintroduced into a plasma deposition source, illustratively at gas inlet25 through mass flow controllers 26. A source of pulsed d.c. or r.f.power 27 is electrically connected to the hollow cathode source forcreating a plasma of high-energy aluminum, nitrogen, and argon speciesin the plasma deposition source 20 in response to applied power, andpreferably pulsed d.c. power between 100 W and 300 W for epitaxialgrowth of zinc-blende AlN. In this embodiment, the power supply is anENI, model RPF-50 set at the standard delay pulse. The source emits alow energy flux 33 of the activated, high-energy aluminum, nitrogen, andargon species. In some embodiments, a negative bias may be applied tothe substrate to accelerate the activated species emerging from thehollow cathode plasma source. In this embodiment, predominantly Al⁺ andN⁺ flow toward the heated substrate and form a film on a front surfacethereof. The process gases are controlled by mass flow control system 30so that the dynamic pressure during deposition, in this particularembodiment, is approximately 5 mTorr. In mass control system 30, gascompositions are monitored and controlled by a differentially-pumped UTImass spectrometer 31 connected through controller 32 for feedbackcontrol of deposition parameters, such as gas flow and partialpressures, substrate temperatures, plasma densities, r.f. target andsubstrate bias, and plasma source induced ion-solid interactions forsynthesis of almost any thin film material.

FIG. 3 is a schematic representation of a magnetically-enhanced hollowcathode plasma deposition source 20 in accordance with the presentinvention. Elements of structure that are common to FIG. 2 are similarlydesignated in this figure. Plasma deposition source 20 comprises acathode 40 defined by a hollow cylindrical chamber 41 having a wall, anupper, or exit, end 42 and a lower end 43. The exterior wall of cathode40 comprises cylindrical magnetic field generating means 44.

In this embodiment, magnetic field generating means 44 comprises anarray of magnets 45 embedded in a cylinder of aluminum 46. In thisspecific example, the array contains 24 magnets: eight magnets aroundthe circumference of the cylinder by three magnets along the length ofthe cylinder. The lengthwise magnets are connected by magnet returns 47.Magnets 45 are, illustratively, rare earth SmCo plated with Ni, and themagnet returns 47 are iron-coated titanium nitride. In this manner, amagnetic field, illustrated by field lines 48, having a strength ofgreater than 1 KG is generated to confine plasma 70 to the hollowcathode.

The interior wall of cathode 40 is source target material 50,specifically MBE grade aluminum in this embodiment, that has been milledto have a thickness that increases from lower end 43 to upper end 42 soas to provide a taper that further confines plasma 70 to the interior ofthe cylindrical chamber. In this specific embodiment, the taper is 3°.

The lower end 43 of cathode 40 is mounted on a conductive cathodemounting block 51. Cathode mounting block 51 includes a water-coolingchannel 49 having a water inlet 52 and a water outlet 53. In thisembodiment, the power supply 27 is integrated into the water coolinglines. An impeller 54 is rotatably mounted in mounting block 51.Impeller 54 may be at ground potential or, alternatively, may be riggedfor bias and electrically isolated from the mounting block. In thisspecific embodiment, impeller 54 is MBE grade aluminum. Impeller 54includes gas inlet 25 for process gases.

The plasma source 20 further comprises an anode 60 which is configuredto surround the cathode at an insulating gap 61. Anode 60 includes a cap62 at the upper end thereof having an opening 63 which is greater thanthe opening 64 at the upper end of the cathode. The opening 64 at theupper end of the cathode emits the low energy flux of activated ions.

The PSMBE system described herein can be used to grow AlN films, as wellas nitride films of other compositions, such as GaN or InN, by varyingthe target and process gas components. The thermodynamically stablewurtzite form of AlN and the metastable zinc-blend form may befabricated selectably in the same PSMBE system by varying the processconditions as illustratively set forth in Table 1.

TABLE 1 Cubic Parameter Hexagonal (wurtzite) A/N (Zinc-Blende) A/NDynamic Gas Pressure 1-10 mTorr 1-10 mTorr Argon Flow 20-40 sccm 20-40sccm Nitrogen Flow 20-40 sccm 2-40 sccm Substrate Temperature 300°C.-900° C. 300° C.-800° C. r.f. Power (to source) 100 W-300 W 0 Pulsedd.c. Power 0 100 W-300 W Acceleration Bias 12 V-15 V (neg.) 0-15 V(neg.) (substrate) Substrates Si(111), Al₂O₃ R-plane), Si(100), 3C-SiC,Al₂O₃ (C-plane), 6H-SiC and MgO(100), or combinations thereof

Of course, in addition to the Group III-V nitride semiconductorsdescribed herein, and their alloys, the PSMBE system can be adapted tofabricate semiconductor devices from other elements, and to formheterostructure devices of varying composition. The system can readilybe adapted, as is known in the art, to include a source of donor oracceptor electrons to form p-n junction devices. For a p-type material,an acceptor source (Group II elements, such as Be, Zn, Cd, and Ca) isincluded during growth whereby the acceptor takes on an electron and isincorporated into the AlN lattice as a negatively charged species. Thesubstrate is bombarded with electrons either by applying a positive biasto the substrate surface or to a metal grid placed directly in front ofthe substrate. For an n-type material, the substrate is bombarded withpositive ions by biasing either the substrate or a grid positioned infront of the substrate negatively so that the donor impuritiesincorporate into the AlN in their charged states (Group IV and VIelements, such as Si, Ge, O, S).

SPECIFIC EXAMPLES

Prior to being loaded in the PSMBE growth chamber described hereinabovewith respect to FIGS. 2 and 3, a Si(100) substrate was de-greasedultrasonically in acetone and methanol and blown dry with nitrogen. Thesubstrate was then etched for 1 minute in 10% (by volume) hydrofluoricacid (HF) and immediately loaded into the PSMBE chamber. The etchingsolution also contained a trace amount of hydrocarbon impurity fromstandard photoresist material. Before deposition, the substrate washeated to approximately 800° C. for one hour which resulted in a (2×1)reconstruction as observed by RHEED.

Using MBE grade aluminum as the target material in themagnetically-enhanced hollow cathode source described in detail withrespect to FIG. 3, a highly purified gas mixture, consisting of 40 sccmAr and 2 sccm N₂, at a dynamic pressure of 5 mTorr, was supplied intothe cathode source. The dynamic pressure of the source gases should becontrolled to the generated plasma stable. The Si(100) substrate washeated to a growth temperature of 650° C. Pulsed d.c. (square wave)power was suppled to the conductive cathode of the source to generateplasma. The d.c. power supply parameters in this specific embodiment areshown below in Table 2

TABLE 2 d.c. power supply parameters Average d.c. Power 100 W Duty CyclePower 25.3% Average Current 1.04 A Peak Voltage −260 V

Several films were deposited on Si(100) substrates in this manner at adeposition rate of approximately 10-15 Å/hr. The impeller in the sourcewas held at ground potential. An acceleration bias of 0 or −12V wasapplied to the substrate. Two films, having a thickness of 800 Å and 960Å, respectively, as measured by piezoelectric resistance, werecharacterized in situ and are discussed hereinbelow. The first film(herein designated “Film 1”) was deposited with an acceleration bias of−12V applied to the substrate. The second film (herein designated “Film2”) was deposited with no bias. The films were not subjected to anypost-growth processing, such as annealing.

Structural Characteristics of the Zinc-blende AlN Films

The films were characterized in situ by a STAIB RHEED system (modelNEK-2035-R) operating at an accelerating voltage of 30 keV during andafter growth. The RHEED patterns indicate that both films have afour-fold symmetry. The orientation of the films with respect to thesubstrate is (100)z-bAlN∥(100)Si and [101]z-bAlN∥[101]Si.

FIG. 4 is a RHEED pattern from Film 1 taken at azimuths of [010], [011],and [031]. Very similar patterns were observed for Film 2. The RHEEDpatterns are spotty and not streaky due to transmission-reflection whichindicates surface roughness. The pattern is formed due to transmissionelectron diffraction from the surface asperities rather than purereflection from the film surface. The presence of spots, as opposed torings, in the RHEED patterns during and after deposition is anindication that the films are single crystalline and that all grains areoriented along the [100] direction.

The films were further characterized by AFM using a Digital InstrumentsNanoscope III in contact mode with standard Si₃N₄ tips. FIG. 5 showsdeflection and height images of the surface of Film 1 taken by AFM.

The height AFM images revealed the presence of regular square pyramidalpits having based edges of up to 2 μm on the surface of both films. Thesquareness of the pyramids indicates that the films are cubic. Itappears that some of these pyramidal pits originate on the AlN/Siinterface. The deflection images show that these pyramids havewell-defined edges and walls. All pyramids are oriented along a certaindirection. From the symmetry of the cubic structure, it can be deducedthat the pyramidal sides are parallel to the {111} planes and thepyramids are oriented along the <110> directions. A height image fromthe surface between the pyramidal pits is shown in the inset andindicates a grain structure having an RMS roughness of approximately 4.4nm. This roughness explains the presence of spots, rather than streaks,in the RHEED patterns and may be an indication that the film was grownin a three-dimensional mode.

It was originally believed that these features were defects formed inorder to reduce the tensile strain in the film due to the large latticemismatch (˜19%) between the zinc-blende film and the silicon substrate.However, reference to FIG. 6 shows a TEM image of a pyramidal pit 9 thatis formed deep into the silicon substrate 6 and is covered evenly withthe zinc-blende AlN film 5. This indicates that a chemical reactionbetween zinc-blende AlN and Si is unlikely. It is known, however, that3C—SiC is deposited on pyramidal pits on Si(100) substrates similar tothose observed in FIG. 5 in the presence of a carbon source. Therefore,it is believed that a layer of 3C—SiC forms on the Si(100) surfaceduring the preheating step. Advantageously, cubic 3C—SiC has a latticeparameter of 4.348 Å which closely matches the lattice parameter ofzinc-blende AlN. An interfacial layer 7 of 3C—Si is shown in FIG. 6.

Thus, in accordance with a preferred method embodiment of the presentinvention, the Si(100) substrate is subjected to a mixture ofhydrocarbons as a source of carbon during processing so that a layer ofcubic 3C—SiC will form on the Si(100) surface during preheating prior todeposition. The high temperature pre-heat (800° C.) decomposeshydrocarbons present on the Si(100) substrate. The source of silicon isthe silicon substrate leading to the formation of pyramidal pits. Theformation of a thin interfacial layer of 3C—SiC provides a cubic matrix,or template, for the growth of cubic zinc-blende AlN. In addition, itacts as a buffer layer that reduces strain in the zinc-blende film.

To confirm the formation of a 3C—SiC film, in this embodiment, at thezinc-blende AlN/Si interface, an AES depth profile was performed onFilms 1 and 2 and the results are shown on FIG. 7. Data were recordedfor Al, N, O, C, and Si in an alternating sputtering mode. A linearleast square algorithm was applied to reduce the perceived noise presentin the depth profile data. Both films (FIGS. 7a and 7 b) clearly showthe presence of a carbon and silicon rich layer at the interface. Thethickness of this layer is estimated to be between 15-25 nm. No tracesof carbon were found at the interface of a sample that was not subjectedto hydrocarbon(s) during the pre-treating procedure (FIG. 7c). While thedeposition conditions for the sample shown in FIG. 7c was identical tothe conditions for Films 1 and 2, this sample did not have a zinc-blendestructure.

Although AES can indicate the presence of silicon and carbon at theinterface between the zinc-blende AlN film and Si(100) substrate, it cannot show whether these elements form SiC, or if SiC is formed, which ofthe more than 250 polytypes of SiC is present at the interface. HighResolution TEM was used to confirm the presence of an intermediate layerat the interface. FIG. 8 is an HRTEM lattice fringe image in the [110]orientation. The intermediate layer is approximately 25-30 Å thick andhas a lighter contrast which is generally indicative of a material withan atomic number smaller than the atomic numbers for the zinc-blende AlNfilm and the Si substrate.

The average {111} interplanar spaces were assessed using a plot-profiletechnique to generate a density profile based on a rectangular selectionfrom the lattice fringe image of FIG. 8 and a row-average plot. The{111} interplanar spacing for the zinc-blende AlN averaged over ten plotprofiles was 2.52 Å while the {111} interplanar spacing for theinterfacial layer was 2.496 Å. These values are in excellentcorrespondence with the theoretical values of 2.53 Å for zinc-blende AlNand 2.515 Å for 3C—SiC. Combined, the results of the AES depth profileand HRTEM confirm that an interfacial layer of 3C—SiC exists in Films 1and 2.

Symmetric (Bragg-Brentano) and rocking KRD scans were performed on thefilms in a Scintag Xl θ-θ diffractometer using Cu K-α radiation and aSi(Li) detector. The samples were supported on zero-background quartzplates. The beam divergence was fixed at 1.4° and scans were collectedin steps of 0.03° (2θ) with either 0.5 or 1 second/point.

FIG. 9 shows the XRD symmetric scans for Film 1 and Film 2. The peakpositions for the zinc-blende AlN layer and silicon substrate arelabeled. Two strong peaks, corresponding to the cubic phase (200)_(C)and (400)_(C) reflection, are prominent. No peaks were observed for the(220)_(C), (002)_(H), or (222)_(C) reflections (the(311)_(C) peak isobscured by the Si (400)peak). A very weak intensity near 2q=36° can bedetected on the XRD pattern of Film 2 and is possibly due to a trace ofeither (002)_(H) or (111)_(C). The measured positions of the (200) cubicpeaks were corrected using the substrate peaks as internal standards.From the corrected positions, the lattice parameter of the cubic phasewas determined to be 4.373±0.01 Å which closely matches the theoreticalcalculation.

Rocking scans about the (200)_(C) peak position are shown for both Film1 and Film 2 on the inset on FIG. 9. These scans display a peakedintensity near q=20.59° with a full width at half maximum (FWHM) of ˜2°.The presence of only (h00) peaks, along with the peaked rocking scans,indicate that the zinc-blende AlN microcrystalline domains are mostly(100) oriented. The small misalignment of about 0.4° for Film 2 could bedue, in part, to sample mounting in the diffractometer.

TEM studies were performed on Film 1 using a JEOL 2010 TransitionElectron Microscope operated at 300 kV. Samples for TEM were preparedusing the standard “sandwich” technique which included grinding,polishing, dimpling, and ion milling at the final stage. All TEM images,in plan view and in cross-section, were taken in bright field imagingconditions.

FIG. 10 shows a cross-sectional view of the AlN/Si(001) interface in the[1{overscore (1)}0] orientation of Film 1. The non-uniform distributionof the intensity in the film reveals regions of microcrystalline domainsand the associated domain boundaries. The image shows a rough andnon-uniform interface between the Si substrate and the AlN film due tothe reaction resulting in the formation of 3C—SiC.

Selected area diffraction (SAD) pattern, taken from numerous regionsacross the AlN/Si(001) interface, repeatedly reveals that the AlN filmis of cubic crystal structure, as shown in FIG. 11. FIG. 11 is plan viewSAD pattern in the [001] orientation which corresponds to a cubiczinc-blende AlN (β-AlN) structure having a lattice parameter of4.383±0.01 Å. The cubic AlN film is single crystalline and epitaxialwith respect to the Si(001) substrate. The very good crystallinity ofthe AlN film is revealed from the shape of the corresponding diffractionspots. Although not perfectly circular, as are the diffraction spotscorresponding to the Si substrate, most of the AlN spots are very welllocalized, with angular distribution within about 1-1.5°.

FIG. 12 is a cross-sectional SAD pattern of the AlN/Si (100) interface.The epitaxial nature of the AlN film is confirmed from the alignmentbetween the spots of the AlN film and the Si substrate: (1{overscore(1)}1)c-AlN∥(1{overscore (1)}1)Si, (2{overscore (2)}0c-AlN∥(2{overscore(2)}0)Si on FIG. 12 and (400)c-AlN∥(400)Si and ({overscore(2)}20c-AlN∥({overscore (2)}20)Si on FIG. 11.

Optical Characterization of the Zinc-blende AlN Films

To date only theoretical calculations of the band gap of zinc-blende AlNhave been reported. Zinc-blende AlN was predicted to be an indirectsemiconductor with a bandgap of approximately 5.11 eV. This value ismore than 1 eV smaller than the bandgap of wurtzite AlN (6.28 eV). Sincethe difference in bandgaps between wurtzite and zinc-blende GaN is only0.2 eV, and equal bandgaps were predicted for wurtzite and zinc-blendeInN, this value was expected to be closer to the bandgap of wurtziteAlN. Moreover, if zinc-blende AlN is indeed an indirect semiconductor,it may be possible to observe an indirect-direct bandgap transition inthe alloys of the zinc-blende forms of AlN with GaN and InN which areboth direct semiconductors.

Film 2 was characterized by spectroscopic ellipsometry using a VariableAngle Spectroscopic Ellipsometer (VASE) over the spectral range 0.73eV-6.25 eV. Some error was introduced in the measurements due to thepresence of the pyramidal pits on the surface of the sample which arelarge enough to scatter light before it can reach the detector.Oscillations in the ellipsometry data were observed to occur at lowerenergies due to interference. A parametric oscillator model was used tofit the optical data for zinc-blende AlN.

The index of refraction (n′) and extinction coefficient (k′) dataobtained by the spectroscopic ellipsometry characterization of Film 2,in the range 1.8-6.25 eV, are presented in FIG. 13. The index ofrefraction of the zinc-blende AN film varied between 2.1 and 2.9 whilethe extinction coefficient was in the range of 0.08 to 0.40. At lowenergy, the data for the extinction coefficient has a constant nonzerovalue, which may be due to the small thickness of the film or thepresence of the pyramids. Ideally, the extinction coefficient will bezero until bandgap absorption takes place. A gradual onset in absorptionis observed to occur at around 5 eV due to interband transitions. Suchgradual onset can be an indication that zinc-blende AlN is an indirectsemiconductor which is in agreement with theoretical predictions. It canalso be due to thickness effects or to the presence of the pyramidalpits on the sample surface. If the extinction coefficient data forzinc-blende AlN are compared to the data from a wurtzite AlN film withapproximately the same thickness, a sharper onset in absorption isobserved for the wurtzite film which is a direct semiconductor. Thiscomparison suggests that zinc-blende AlN, fabricated in accordance withthe present invention, is indeed an indirect semiconductor.

When the conduction band minium occurs at the same point in k space asthe valence band maximum, the semiconductor is direct. If the minium ofthe conduction band occurs at a different point in k space than themaximum of the valence band, then the semiconductor is indirect. Basedon the assumption that zinc-blende AlN has an indirect bandgap, thebandgap of zinc-blende AlN was estimated, from the extinctioncoefficient data. If the phonon energies compared to the totaltransition energy is ignored, absorption from interband transitions hasthe form:

α_(g)(E)=C(E−E _(g))¹ ^(_(ν)) ⁻¹,  Eq. (1)

where 1=2 if the vertical optical transition at k=0 is allowed and 1=3if the transition at k=0 is forbidden, E=hν is the photon energy, andE_(g) is the bandgap. The extinction coefficient is related to theabsorption by:

k′=αλ/4π  Eq.(2)

By re-plotting the extinction coefficient data as (Eα)^(½) as a functionof the photon energy at 1=2, the lowest indirect bandgap can be obtainedfrom a linear fit to the appropriate range of the curve. FIG. 14 showssuch a plot. The data exhibit two ranges of liner dependence. The one atlower energies may be related to the presence of pyramids on the samplesurface or a thickness effect. The linear range at higher energies isrelated to interband transitions, as shown in Eq. (1). The bandgap wouldbe approximately the x-axes coordinate of the intersection point of twotangents as shown in FIG. 14. The first is the tangent to that part ofthe plot which corresponds to surface roughness or thickness effect andthe second tangent follows the part of the plot which corresponds toabsorption from the bandgap. In this fashion, the bandgap of zinc-blendeAlN is approximately calculated to be at 5.34 eV. This value is inreasonable agreement with the theoretical value of 5.11 eV previouslyreported and confirms that the bandgap of zinc-blende AlN is on theorder of 1 eV smaller than the bandgap of wurtzite AlN.

The material of the present invention has many uses, including but notlimited to, use as a substrate and as a template for growing cubiccrystals with conventional systems, such as a chemical vapor depositionsystem. Although this wide bandgap semiconductor is somewhat resistantto defects or imperfections in the material, it is believed that lessdefects will occur in an epilayer grown on a cubic substrate.

Among the anticipated uses of cubic zinc-blende AlN are ultravioletphotonic detectors and surface acoustic wave sensors. It can also beused as a dielectric coating for enhancement of Kerr rotation in opticalstorage media and as an insulator for silicon carbide-based devices.Furthermore, the wurtzite polytypes of InN, GaN, and AlN form acontinuous alloy system with direct band gaps ranging from 1.9 eV forInN to 6.2 eV for AlN. Thus, AlN in combination with other nitridescould potentially be fabricated into optical devices that are active atoptical wavelengths ranging from red well into the UV. Moreover, sincethe cubic structure is expected to be easy to dope, the material will beuseful for p-n junction devices.

The following is a list of exemplary devices, and is in no way intendedto be limiting: surface or other acoustic wave devices for chemical gasand biological detectors, heterojunction field effect transistors,ultraviolet photonic sensors, pyroelectric and piezoelectric-basedsensors, field effect transistors and bipolar junction transistors,microwave and millimeter wave detectors, schottky and p-n junctiondiodes, metal-insulator-semiconductor capacitors, microwave field effecttransistors, radiation detectors and x-ray/radiation imagers, flat paneldisplay electron emitters, and thermal (electrically insulating) heatsinks for electronic packaging.

In a specific illustrative device embodiment, shown in FIG. 15, aninterdigitated transducer SAW device 70 comprises a substrate 71,illustratively Si(100), a buffer layer 77 of 3C—SiC, an epitaxial film72 of single crystal, zinc-blende AlN oriented (100) on the substrate.The zinc-blende AlN film 72 has a patterned film of two interdigitatedelectrodes 73 and 73′ deposited on it by standard photolithographictechniques. A signal generating source 74 is connected across one end ofthe electrodes 73 and a signal detector 75 is connected across the other(73′). When an alternating voltage is applied to the electrodes, asurface acoustic wave is generated that travels along the device in thedirection of arrow 76 and is detected by signal detector 75.

Although the invention has been described in terms of specificembodiments and applications, persons skilled in the art can, in lightof this teaching, generate additional embodiments without exceeding thescope or departing from the spirit of the claimed invention.Accordingly, it is to be understood that the drawing and description inthis disclosure are proffered to facilitate comprehension of theinvention, and should not be construed to limit the scope thereof.

What is claimed is:
 1. A semiconductor device comprising a film ofdevice-quality, single crystal cubic zinc-blende AlN having a thicknessof at least 800 Å.
 2. The film of claim 1 wherein the cubic zinc-blendeAlN has a lattice parameter of approximately 4.373 Å.
 3. A semiconductordevice comprising: a substrate; and a layer of single crystal epitaxialzinc-blende AlN on a face surface of the substrate, the layer ofepitaxial zinc-blende AlN having a thickness of at least 800 Å.
 4. Thesemiconductor device of claim 3 wherein the substrate has cubic symmetryon its growth surface.
 5. The semiconductor device of claim 4 whereinthe substrate is selected from the group consisting of Si(100),MgO(100), 3C—SiC and combinations thereof.
 6. The semiconductor deviceof claim 4 wherein the cubic substrate is Si(100) with a surface layerof 3C—SiC.
 7. A semiconductor device comprising: a Si(100) substrate; alayer of 3C—SiC on a first surface of the Si(100) substrate; and a filmof single crystal epitaxial zinc-blende AlN on the layer of 3C—SiC,oriented epitaxial to the (100) plane.
 8. The semiconductor device ofclaim 7 wherein the zinc-blende AlN has a thickness of at least 800 Å.9. The semiconductor device of claim 8 wherein the zinc-blende AlN has athickness in the range of at least about 1000 Å to 2000 Å.
 10. A surfaceacoustic wave device comprising: a substrate; a single crystal epitaxiallayer of cubic zinc-blende AlN having a thickness of at least 800 Ådeposited on the substrate; and electrodes on the surface of theepitaxial layer of cubic zinc-blende AlN for converting an electricalsignal to a surface acoustic wave and for converting the generatedsurface acoustic wave into an electrical signal.
 11. The surfaceacoustic wave device of claim 10 wherein the substrate has a cubicsymmetry on its growth surface.
 12. The surface acoustic wave device ofclaim 11 wherein the substrate is Si(100) with a layer of cubic 3C—SiCthereon.
 13. The semiconductor device of claim 9 wherein the zinc-blendeAlN has a thickness in the range of at least about 1000 Å to 2000 Å.